Method for processing dispersion strengthened metals



3,979,439 Patentecl Dec. 25, 1962 3,076,439 METHQD FOR PROCESSINGDISPERSION STRENGTHENED METALS Nicholas J. Grant, Winchester, and KlausM. Zwilsky,

Watertown, Mass, assigners to New England Materiais Laboratory, Inc,Medford, Mass, a corporation of Massachusetts No Drawing. Filed Mar. 15,1960, Ser. No. 15,036 11 Claims. (Cl. 75206) This invention relates to amethod of processing dispersion strengthened or hardened metals and, inparticular to a method of hot working said metals while inhibiting thedissipation of stored energy arising from the strain deformation of suchmetals.

Dispersion strengthening of alloys consists of dispersing a finelydivided non-metallic phase throughout a metal matrix and then storingenergy of deformation in the structure. Several techniques have beenemployed to obtain the desired dispersion and among these, mechanicalmixing and internal oxidation have thus far shown the most promise.Mechanical mixing is the fastest and least expensive way to obtaindispersion strengthened materials. Internal oxidation, by comparison, isa more expensive and time consuming process to operate on a larger scaleand involves the use of fine alloy powders not readily available. toryoxide-forming metal, e.g. aluminum, is alloyed with a matrix metal, suchas copper, and the alloy in particulate form subjected to an oxidationtreatment to oxidize selectively aluminum to a dispersion of A1 followedby consolidation of the particles to a wrought shape. In terms ofinterparticle spacing, the latter technique is usually capable ofproviding wrought shapes exhibiting markedly improved results forsmaller amounts of the non-metallic phase than wrought shapes producedfrom mechanically mixed powder.

From an economical viewpoint, mechanical mixing of metal and refractoryoxide is preferred for the processing of common metals such as iron-basealloys.

Investigations of certain metal and refractory oxide systems haveindicated that the size of the metal powder, the size of the oxidepowder and the volume percent oxide are all important variables inachieving optimum properties.

With regard to dispersion strengthened alloys, stress rupture data haveshown that such alloys are generally characterized by extreme flatnessof slope on a log stress versus log rupture life plot. Because of thischaracteristic, the superiority of such alloys over conventionalmaterials becomes markedly effective as time and/or temperature areincreased. This is due to the fact that conventional materials, such asage hardenable alloys, are subject to over aging or solution of a secondphase as affected by time and temperature, whereby a rapid drop-off inproperties occurs. On the other hand, oxide dispersion strengthenedmaterials have been known to be generally stable up to below the meltingpoint of the base or matrix material.

Strengthening in these alloys has been postulated to arise from twosources. One is the effect of the dispersion, which imparts strength tothe matrix metal even in the annealed condition. The other efiFectarises from the introduction of the strain energy during the workingprocess, tag. the extrusion process, with the consequent retention ofthis energy by the alloy by virtue of the presence of the dispersedphase.

We have found, however, that not all matrix metals behave the same wayand that in some instances, unless due care is taken during straindeformation in producing a wrought shape from the dispersionstrengthened alloy, at least some of the stored energy is dissipated. We

In this process, a refracfound this to be particularly true for iron andalloys. v

It is an object of this invention to provide a method of deformingdispersion strengthened metals whereby stored energy is imparted theretowith the minimum of energy dissipation while enhancing the strengthproperties of said metal.

Another object is to provide a method of processing dispersionstrengthened iron and iron alloys with the aim of producingmaterial ofenhanced strength properties, such as improved resistance to creep.

It is a further object to provide a method for processing a dispersionstrengthened metal wherein the matrix metal is characterized by a phasetransformation at an elevated temperature, and wherein the deformingtemperature employed in processing said metal is determined according tothe phase transformation of said metal.

An additional object is to provide a method of processing a matrixmetal, such as iron or an iron-base alloy, having dispersed therethroughfine particles of a refractory oxide wherein the processing temperatureis determined according to the temperature of phase transformation ofthe matrix metal and the temperature of crystallographic transformationof the refractory oxide.

These and other objects will more clearly appear from the followingdisclosure.

In producing a dispersion strengthened metal, such as iron, finelydivided iron is mixed with up to about 15 W0 (volume percent) of arefractory oxide, such as A1 0 the particle size of the iron beingpreferably below 20 microns and generally below 10 microns. We preferfor optimum results to use particle sizes below 5 microns. We find itimportant that the refractory oxide have a particle size less than thatof the matrix metal, for example from about 30 to 250 times smaller,preferably from about 30 to times smaller, so as to obtain adequatedistribution of the oxide and control the interparticle spacing betweenthe oxide particles. The metal powder is mixed with the refractory oxidepreferably in the dry state by means of a high speed dry blender, or aball mill, for a suitable time to effect uniform mixing, e.g. rangingfrom about 15 to 60 minutes in the blendor and 1 to 24 hours for theball mill. Upon completion of the mixing, the mixture is subjected to areduction treatment with hydrogen to reduce any iron oxide present at anadequate reducing temperature, e.g. from 700 to 900 F., and thenhydrostatically pressed to the desired iron-base shape, followed bysintering in hydrogen to form a compact capable of being handled. Wefind that the sintering temperature should preferably not exceed thetransformation temperature of the matrix metal. Thus, in the case ofiron, we prefer the sintering temperature be maintained in the alphairon region, for example at 1525 F. This temperature is also importantwhere gamma alumina is used as the dispersoid as it is below the gammato alpha transformation temperature. Thus, severe agglomeration of thepowders is substantially avoided. The sintered compact is thereaftervacuum packed in a can and the whole extruded at an elevated temperatureat an extrusion ratio sufficient to achieve maximum density, e-.g.ranging from about 10 to 1 to 28 to l.

We have found that dispersion strengthened iron or iron-base alloysproduced by the foregoing powder metallurgy or similar techniquegenerally exhibit markedly improved strength properties, provided thatthe deformation employed in producing a dispersion strengthened metalshape of maximum density is carried out while the material is in theferritic or alpha condition. In other words, we have found that, inproducing a dispersion strengthened metal product from a matrix metalpowder characterized by a phase transformation at an elevatedtemperature and having mixed therewith fine particles of a non-metallicdispersoid, such as alumina, the consolidation of the mixture should beconducted at an elevated temperature below the temperature at whichphase transformation occurs inorder to insure optimum strengthproperties. Our tests have indicated that if the Working or. deformingis conducted above the phase transformation temperature of the matrixmetal, the stored energy is dissipated as the worked material cools downthrough the, phase transformation range of the material and convertsfrom one crystallographic structure to another; for example, as in ironfrom austenite to ferrite.

In order to obtain a better understanding of the invention, thefollowing examples are given:

Iron powder of about 3 microns average particle size was mixed withvarying amounts of up to about 10 v/o of alumina having an averageparticle size of about 0.027 micron, the ratio of particle size of ironto aluminum being a little more than .100 to l.

The alloys were prepared by dry mixing the iron matrix metal'powder andthe aluminum oxide in lots of about 500 grams, the mixing beingconducted in a high speed blendor, e.g. a Waring Blendor, at a speed ofabout 15,- 000' revolutions per minute; The mixing of each lot wascarried out for about 5 minutes and then further mixed by spatulation ona sheet of clean paper for a few minutes, the procedure including theblendor and subsequent spatulation being repeated about 4 times.

The blended powder batches were thereafter subjected to a reducingtreatment in dry hydrogen for a minimum of five hours at a temperatureof about 800 F. to insure clean particle surfaces for subsequentconsolidation of the mixture into wrought shapes. Each batch of themixed powders was introduced into a rubber tube Supported within aperforated steel canister about two inches in diameter, one end of therubber tube being rubber stoppered at the start. After the powder wasintroduced, a second rubber stopper having in communication therewith ahypodermic needle was inserted, a vacuum connection being made throughthe needle to remove the air from within powder mass. After completionof evacuation, the needle was removed and, the canister assemblysubjected to hydrostatic pressure at about 30,000 p.s.i. to yieldcompacts about 1.4 inches in diameter and 3 inches long.

Using correspondingly larger amounts of powder, compacts have beenprepared by the same technique up to 3 inches in diameter and 6 incheslong. Depending upon the size of presses, billets of up to 20 inchesdiameter are envisaged.

The compacts produced as aforementioned were then subjected to sinteringin dry hydrogen for a minimum of -hours at 1525 F. After that they wereeach canned by insertion in a mild steel can and welded vacuum tightfollowed .by extrusion at an elevated temperature. The extrusion ratiowas about 16 to 1.

The following alloys were produced:

Table I V01. V01. Extru- Al loy No. percent percent 2 sion F81 A1203 tem? 1 3 micron iron powder. 2 0.027 micron A1203 powder.

Alloys? and 4 of the same identical composition were extruded atdifierent temperatures, one at 1550 F. (in the ferritic range), theother at l9.00 F. (in the austenitic range). V

The alloys were then subjected to tensile tests at room worked No. 4.

Table II Y.S. Alloy No. Temp., (p.s.i.) Ultimate Percent F. (0.2%(p.s.i.) elong,

offset) R.T. 66, 230 8'), 800 20 1, 270 13, 700 19,100 35 .T. 68, 00084, 500 15 1, 2l0 16, 200 21, 25 RIP. 63, 900 98,000 9 1, 230 21, 30028, 400 14 1, 400 13, 900 18, 300 9 4 R.T. 69, 800 88, 400 8 1, 230 10,800 16, 200 20 1, 400 6, 200 9, 000 16 5 R.T. 104, 300 113, 000 6 1, 23026, 830 33, 300 10 Table HI Vol. 10') hr. 1200 F. Alloy No. percentrupture life, percent A1203 stress elong. (p.s.i.)

4 11,000 6 e 13, 800 s a g is, 000 2 8 6, 500 3 10 24, 000 2 Forcomparison purposes, an extrusion was made from the same iron powderfree from the presence of A1 0 This material designated as F-l,exhibited, in the asextruded condition, a yield strength at roomtemperature of about 25,000 p.s.i. and at 1000 F. of about 6,600 p.s.i.while exhibiting a 100 hour rupture life (at 1200 F.) under a stress of2,600 p.s.i. (12% elongation).

The results indicate that the properties of iron are markedly improvedby the presence of A1 0 as a dispersion strengthener. The improvement isparticularly noticeable with respect to strength properties at elevatedtemperatures. In this connection, reference is made to the 100 hourrupture properties which show an increase in rupture life stress oversimilarly prepared pure iron of about 4 to 9 times for alumina contentsranging from 4 to 10 v/o.

Our investigations have indicated that in order to achieve the markedlyimproved results, it is important that the fabrication of the materialcontaining the dispersion strengthener be carried out below the phasetransformation temperature of the matrix metal. This is illustrated bythe results obtained for alloy Nos. 3 and 4, both of which contained 8vol. percent A1 0 with the exception that No. 4 was extruded at 1900 F.,i.e.

in the austenitic region of the matrix metal iron, while No. 3 wasextruded at 1550 F., i.e. in the ferritic region of the matrix metal.Referring to Table II, it will be noted that No. 3 exhibited a yieldstrength at 1200 F. of 21,300 p.s.i. as compared to the value of 10,800p.s.i. (about half) obtained for No. 4 which was subjected to austeniticextrusion. Similarly, at 1400 F. No. 3 exhibited a yield strength of13,900 p.s.i. as against the much lower value of 6,200 p.s.i. for theaustenitically V The same trendwas indicated with respect to the 100hour rupture'life at. 1200" F. wherein the ferritically worked No. 3alloy exhibited a stress of 18,000 p.s.i. as against the austeniticallyworked No. 4 which exhibited a stress of 6,500 p.s.i., almost one-thirdthe former value. It is thus apparent that unless the working isconducted below the transformation temperature of the matrix metal, theelevated temperature properties of the wrought metal are adverselyaffected.

Tests in which MgO was used as the dispersion strengthener forferritically worked alloys were alsocon- -changes at elevatedtemperatures.

ducted. Using the same particle size iron powder (3 microns) and MgOwith its particle size ranging from about 0.05 to 0.1 micron, roomtemperature yield strengths-of 84,500; 102,000 and 129,000 p.s.i. wereobtained for iron alloys containing 4, 6 and 10 v/o, re-

spectively, ranging from over 3 to over 5 times the value of thatobtained for pure iron. Likewise, yield strength values at 1000 F. wereobtained ranging from about 22,800 to 30,600 p.s.i., as, against 6,600p.s.i. for pure iron similarly prepared. As for 100 hour rupture life atworking of the MgO-containing alloys was likewise es- ;sential inobtaining optimum properties.

Observations in other dispersion strengthening systems ihave indicatedthat in order to obtain optimum properties -consistently, considerationshould be given to the tendency of the dispersoid to undergocrystallographic Microstructural examinations of some dispersionhardened alloys of other systems have indicated that when an alloy isworked at a temperature above the crystallographic phase transformationof a dispersoid, for example alumina, agglomera- -tion of the dispersoidhas been observed to occur as the dispersoid transforms from onecrystallographic phase to another, thus partially destroying theinherently fine dispersion of the dispersoid with a consequent fallingoff in strength properties. In the case of alumina employed Thistransformation starts in excess'of 1500 F. and is believed to becomplete at about 1850 F. Therefore, we prefer,

where a dispersoid is employed which crystallographically transforms atan'elevated temperature, that the temperature of consolidation of thealloy be also maintained below this transformation temperature as wellas below the phase transformation of the matrix metal.

While the present invention has been described with respect to itsapplication to iron, it is also applicable to iron-base matrix alloyssuch as carbon steels, certain of the ferritic stainless steels subjectto gamma or other transformation, certain of the iron-base alloy steelsand the like. Examples of iron-base alloys would be those single phasebinary alloys which have a second metal which does not oxidize readilysuch as binary alloys of Fe-Cr (up to 20% Cr), FeW (up to 4% W), FeMo(up to 4%), FeNi (up to 10% Ni), FeSi (up to Si), etc. It will beappreciated that some of these alloys may have an alpha or gamma phasestable up to very high temperature in which case they may not presentany problems. However, these ranges may include small amount of otherelements such as carbon, manganese, or other strong austenite formerssufiicient to effect transformation. Ternary and more complex alloys,such as FeCr-Mo, FeNiW and the like would be ineluded where they aresubject to phase transformation at elevated temperatures. The termiron-base alloys is meant to include solid solution alloys containing atleast 50% Fe and, preferably, at least 65% Fe.

Other matrix metals and alloys to which the invention is applicable arecobalt (transforms between 800 to 900 F.), and such cobalt alloys asCo-Mo, CoNi-W, CoCrW, CoCrMo, etc., reference being made to availablephase diagrams in determining adequate fabrication temperatures freefrom phase transformations. Still others are titanium (a phasetransformation at about 1615 F.), zirconium (a phase transformation atabout 1470 F.), uranium (a phase transformation at about 1115 F.) aswell as other metals and alloys.

For the purposes of this invention, it is preferred that the matrixmetal have a melting point above 1800" F. Where such metals are employedas the matrix, as described hereinbefore the metal in particulate formis .rnixedwith a refractory oxide powder, shaped and converted into aform capable of being easily handled during deformation working, and thewhole but worked at an elevated temperature below the phasetransformation temperature of the matrix metal and also preferably beloWthe crystallographic transformation temperature of the dispersedrefractory oxide.

Examples of refractory oxides which may be used in producing dispersionhardened alloys in accordance with the invention are SiO A1 0 MgO, CaO,BaO, SrO, BeO, ZrO TiO Th0;; and oxides of the rare earth metal group,such as oxides of cerium, lanthanum, neodymium, etc. Such refractoryoxides are characterized as having a melting point above the meltingpoint of the matrix metal, for example above 2700 F. and generally aboveIn addition, these oxides are characterized by a negative free energy offormation at about 25 C. of over 90,000 calories per gram atom ofoxygen. For example SiO has .a negative free energy of formation at 25"C. of about 96,200, A1 0 of about 125,590, MgO of about 136,130, BeO ofabout 139,000, etc. Where such oxides exhibit crystallographictransformation at elevated temperatures, such as SiO A1 0 ZrO TiO andothers, then we prefer, depending on the crystallographic phase present,that the working temperatures employed be maintained below thetransformation temperature of the oxide phase to inhibit agglomerationthereof to larger sizes. For examples, where A1 0 is used as thedispersoid and it is in the gamma form, then the temperature offabrication of the alloy should not exceed the crystallographictransformation temperature for gamma alumina which falls within range ofabout 1500 F. to 1920" F., otherwise agglomerations of the aluminaparticles occurs. However, where the alumina used is already in thestable alpha form, the temperature of fabrication makes no difference inso far as the alpha alumina is concerned as it is in the stable hightemperature form. SiO exhibits several transformations, ZrO has amonoclinic form which is stable to about 1830 F.

When using oxides of the foregoing types as dispersoids, as statedhereinabove they may range from about 1 to 15 We and preferably fromabout 3 to 12 v/o. Likewise, as stated hereinbefore, the particle sizeof the oxide should be 30 to 250 times smaller than the matrix metalpowder, preferably 30 to times smaller, and also preferably range insize from about 0.01 to 0.1 micron for particle size of the matrix metalpowder ranging up to about 20 microns or over the range of 1 to 20microns.

The invention enables the production of high strength Wrought metalproducts characterized by structural stability and. by the ability toretain high strength properties at elevated temperatures. Examples ofmetal structures are heat exchangers, turbine buckets of iron or ironalloys and other metals for use in steam and gas turbines, furnacestructures where resistance to creep at elevated temperatures is animportant consideration, boiler tubing for carrying super-heated steamand many other applications too numerous to mention.

Although the present invention has been described in conjunction withpreferred embodiments, it is to be understood that modifications andvariations may be re- 7 sorted without departing from the spirit andscope of the invention as those skilled in the art will readilyunderstand. Such modifications and variations are considered to bewithin the purview and scope of the intention and the appended claims.

What is claimed is:

1. A method of producing a dispersion strengthened wrought metal productfrom a matrix metal selected from the group consisting of iron, cobalt,iron-base and cobaltbase metals of melting point above 1800 F.characterized by a phase transformation temperature at an elevated hotworking temperature and from an insoluble refractory oxide dispersionhardener characterized by a negative free energy of formation at about25 C. of at least about 90,000 calories per gram atom'of oxygen whichcomprises, providing a powder mixture comprising said matrix metal and asubstantially uniform dispersion of said shape atan'elevatedtemperature'not exceeding the temperature at which the phasetransformation of said matrix metal begins, whereby dissipation ofstored energy is greatly inhibited.

2. The method of claim 1, wherein the amount of dispersed refractoryoxide particles ranges from about 1 to 15 volume percent of'the volumeof the metal product.

3. The method of claim 2, wherein the particle size of the oxide rangesfrom'about.0.01 to 0.1 micron.

4. The method of claim 1, wherein the compact is hot deformed byextrusion.

5. A method of producing'a dispersion'strengthened wrought metal productfrom an iron-base matrix metal of melting point above 1800 F.characterized by a ferritic to austenitic transformation temperature atan elevated hot working temperature and froman insoluble refractoryoxide dispersionhardener characterized by a negative free energyof-formation at about 25 C. of at least about 90,000 calories pergramatom ofoxygen which comprises, providing a powder mixture comprisingsaid matrix metal and a substantially uniform dispersion of saidrefractory oxide powder, forming a compact of said mixture,hot deformingsaidcompact to" a wrought metal shape at an elevated temperature not"exceeding the temperatureat which the phase transformation of saidmatrix metal begins, whereby dissipation of stored'energy is greatlyinhibited.

6. The method of claim 5, wherein the iron-base metal is comprisedsubstantially of iron.

7. The method of claim 5, wherein the compact is hot deformed byextrusion.

8. A method of producing a dispersion strengthened wrought metal productfrom an iron-base matrix metal of melting point above 1800 F. comprisedsubstantially of iron characterized by a ferritic to austenitictransformation temperature at anelevated hot working temperatureand'frorn an insoluble refractory oxide dispersion hardenercharacterized by a negative free energy of formation at about 25 C. ofat least about 90,000 calories per gram'atom of oxygen which comprises,providing-a powder mixture of said matrix metal of average particle sizeranging up to about 20 microns with about 1 to 15 volume percent of asubstantially uniform dispersion of said refractory oxide powder ofaverage particle size ranging from about 0.01 to 0.1 micron, formingacompact of said mixture, and hot deforming saidcompact to "a wroughtmetal shape at an elevated temperature not exceeding the temperature atwhich the phase transformation of said matrix metal begins, wherebydissipation of stored energy is greatly inhibited.

9. The method of claim 8, wherein the compact is-hot deformed byextrusion.

.10. The method of claim 8, wherein the refractory oxide is alumina.

11. The method of claim .10, wherein the amount of alumina ranges'fromabout 3 to 12 volume percent.

References Cited in the file of this patent ,UNITED STATES PATENTS2,823,988 .Grantet'alL Feb. 18, 1958 2 ,855,659 Thomson Oct..l4, 19582,894,838 Gregory July 14, 1959 OTHER REFERENCES Transactions AIME, partin Journal of. Metals," February 1954, pages 247-249.

1. A METHOD OF PRODUCING A DISPERSION STRENGHENED WROUGHT METAL PRODUCTFROM A MATRIX METAL SELECTED FROM THE GROUP CONSISTING OF IRON, COLBALT,IRON-BASE AND COLBALTBASE METALS OF MELTING POINT ABOVE 1800*F.CHARACTERIZED BY A PHASE TRANSFORMATION TEMPERATURE AT AN ELEVATED HOTWORKING TEMPERATURE AND FROM AN INSOLUBLE REFRACTORY OXIDE DISPERSIONHARDENER CHARACTERIZED BY A NEGATIVE FREE ENERGY OF FORMATION AT ABOUT25*C. OF AT LEAST ABOUT 90,000 CALORIES PER GRAM ATOM OF OXYGEN WHICHCOMPRISES, PROVIDING A POWDER MIXTURE COMPRISING SAID MATRIXMETAL AND ASUBSTANTIALLY UNIFORM DISPERSION OF SAID REFRACTORY OXIDE POWDER,FORMING A COMPACT OF SAID MIXTURE, AND HOT DEFORMING SAID COMPACT TO AWROUGHT METAL SHAPED AT AN ELEVATED TEMPERATURE NOT EXCEEDING THETEMPERATURE AT WHICH THE PHASE TRANSFORMATION OF SAID MATRIX METALBEGINS, WHEREBY DISSIPATION OF STORED ENERGY IS GREATLY INHIBITED.